Bainitic ferrous alloy and method

ABSTRACT

This invention relates to an as-worked bainitic ferrous alloy and to a novel method of processing same to obtain optimum strength and toughness. More particularly, this invention is directed to the hot working cycle of a ferrous alloy characterized by an I-T Diagram or &#39;&#39;&#39;&#39;S&#39;&#39;&#39;&#39; Curve having a double nose or a pearlite transformation knee of the beginning curve above a broad bainitic bay region. Such an alloys is heated to an austenitizing temperature of about 1,500* to 2,200* F., and subjected to a plurality of working operations at successively lower temperatures, where the final working operation is conducted after the beginning of the austenite transformation to bainite and before the complete transformation thereof.

United States Patent 1191 Bramfitt et al.

1451 Sept. 23, 1975 BAINITIC FERROUS ALLOY AND METHOD [73] Assignee: Bethlehem Steel Corporation,

Bethlehem, Pa.

22 Filed: Feb. 27, 1974 21 Appl. No.: 446,550

Related US. Application Data [62] Division of Ser. No. 316,962, Dec. 20, l972.

Primary Examiner-W. Stallard Attorney, Agent, or FirmJoseph J. OKeefe; William B. Noll [57] ABSTRACT This invention relates to an as-worked bainitic ferrous alloy and to a novel method of processing same to obtain optimum strength and toughness. More particularly, this invention is directed to the hot working cycle of a ferrous alloy characterized by an l-T Diagram or S Curve having a double nose or a pearlite transformation knee of the beginning curve above a broad bainitic bay region. Such an alloys is heated to an austenitizing temperature of about l,500 to 2,200 F., and subjected to a plurality of working operations at successively lower temperatures, where the final working operation is conducted after the beginning of the austenite transformation to bainite and before the' complete transformation thereof.

4 Claims, 1 Drawing Figure 521 u.s.c1. 148/36 51] 1111.131. ..c22c 38/12 [58] Field 61 Search 148/36, 12 F [56] References Cited UNITED STATES PATENTS 3.303.061 2/1967 Wilson 148/36 3348981 10/1967 Goda 6161. 148/36 3,463,677 8/1969 Nakamura 6161.. 148/36 3,753,796 8/1973 MellOy et al 148/36 Standard H Roma llfro TIME TEMPERATURE U S Patent sep 23,1975

Standard H01 Rolled h'OI/ed Ho INVENTION TIME BAINIITIC FERROUS ALLOY AND IVIETHOD This is a division of application Ser. No. 316,962, filed Dec. 20, 1972.

BACKGROUND OF THE INVENTION The invention, to be described in detail hereinafter, is directed to an as-worked bainitic ferrous alloy, and to a process including the thermomechanical treatment thereof. In general terms, bainite has been metallurgically defined as one of the transformation products of austenite. Upon cooling the austenite, transformation to bainite occurs over the temperature range of about 1,000" to 450 F. The microstructure differs from pearlite, a high temperature austenite transformation product, in that it is acicular in nature. The last generally recognized transformation product is martensite. Transformation to martensite must be directly from austenite, i.e. without prior transformation to ferrite, pearlite or bainite. Unlike the transformation to pearlite or bainite, it is not time dependent, but occurs almost instantly upon cooling.

Over the recent years, the times, temperatures and rates of transformation of various steels have been determined and plotted the result being a series of diagrams known as isothermal transformation diagrams, 'ITT diagrams, or S-curves. One current publication offering l-T and C-T or continuous cooling diagrams for standard steels, as well as modified versions thereof, is the Atlas of Isothermal Transformation Diagrams by U.S. Steel, 1963. It will be appreciated that the temperature and range noted above is not limiting but merely illustrative of atypical steel. A quick review of an l-T diagram for a given steel will give a closer reading of said steel. Nevertheless, while they are close, such diagrams are not precise as to exact times and temperatures.

For example, the I-T diagrams were developed by studying the transformation behavior of a steel at a series of temperatures below the A critical temperature, by quenching small samples to the desired temperature in a liquid bath, allowing them to transform isother mally and following the progress of the transformation metallographically. From this, the curves could be drawn. Very generally, this procedure included heating, quenching, holding, and cooling. It has since been determined that processing variations such as working or rolling affected the actual time and temperature of transformation, particularly the pearlite transformation. More specifically it was discovered that the inclusiori of an austenite working step changed the transformation kinetics of the steel by shifting the curve to a higher temperature. This phenomena has been confirmed by Y. E. Smith in an article in Metallurgical Transactions, V. 2, June 1971. Nevertheless, with the practices have been developed with the most well known of these being ausforming". This is identified and shown in the FIGURE, where the practice is schematically illustrated superimposed on a typical I-T curve of a bainitic ferrous alloy. Isoforming is anotherpractice where the ferrous alloy is worked as it isothermally transforms to pearlite. This, along with standard and controlled hot rolling, is illustrated in the FIGURE in the manner of ausforming.

Variations to certain of these practices have been developed as evidenced by the following US. Pat. Nos. 2,717,846 3,215,565 2,240,634 3,303,061 3,340,102 3,444,008 3,453,152 and 3,645,801. While some of said patents treat alloys susceptible to processing by the method herein, none of them individually or in combination teach the present concept of subjecting the alloy in an austenitized condition to a plurality of working operations, preferably at least five, where the final working is concluded after the initiation of the austenite transformation to bainite and before the complete transformation thereof.

SUMMARY OF THE INVENTION This invention is directed to an as-worked bainitic ferrous alloy and to a procedure involving the thermomechanical treatment of the said alloy. More particularly it relates to a process of heating said alloy to an austenitizing temperature between about 1,500 to 2,200 F. and subjecting the alloy to a plurality of working operations as it cools from said austenitizing' temperature. While the final working should be conducted during the transformation of the austenite to bainite, the initial working should begin while the alloy is above l,500 F. By this procedure, an optimum combination of strength and toughness, in the as-worked condition, is produced. Metallographically, the processed alloy is characterized by fine grains and cells, with very fine uniformly dispersed carbides.

BRIEF DESCRIPTION OF THE DRAWING The FIGURE represents a typical I-T curve of a bainitic alloy susceptible to treatment by the process herein, with a number of schematic representations of thermomechanical treatment cycles superimposed thereover, comparing the present process to that of four prior art processes.

DETAILED DESCRIPTION OF PREFERRED EMBODIMENT In the preferred practice of this invention, a bainitic ferrous alloy characterized by an isothermal transformation diagram having the pearlite transformation knee of the beginning curve laterally displaced to the right of the lower temperature bainitic transformation knee, is heated to an austenitizing temperature between about 1,500 to 2,200 F., preferably at least about 1,600 E, and subjected to at least two working steps, preferably at least five, to reduce the cross section thereof, with the final of said steps occurring within the bainite transformation region.

Without intending to unduly limit the invention, a number of ferrous base alloys having the foregoing characteristic are listed with nominal compositions in Table I, by weight percent:

TABLE I Baintic Ferrous Alloy C Mn Mo Cr Ni V Si Al B Cu A .39 .56 .74 3.53 B .23 .82 .53 1.22 .22 C .40 .78 .53 1.25 .22 D .33 .84 1.07 1.05 .26 E .27 .84 .90 .73 .60 .l 1 F .25 .88 .88 .73 .59 .23 G .6 .45 1.52 3.33 H .55 .60 .19 1.03 .36 J .42 .78 .33 .80 1.79 K .62 164 .32 .60 1.79 .67 L .59 .89 .22 .64 .53 M .57 .82 .26 1.07 1.16 N .55 .83 .48 1.01 1.15 P .51 .73 .45 .99 2.75 Q .41 .57 .36 1.57 1.26 R .46 .79 .18 .77 .91 .0021 S .15 .92 .46 .50 .88 .06 .0031 .32 T .39 .89 .50 .95 .88 .03 .48 .002 U .44 .79 .54 2.10 .06 1.63 V .26 .77 1.00 .95 .04 .25 W .035 .79 .97 .95 .02 .24

Each of said bainitic alloys A-W are characterized by a double nose I-T Diagram where the upper or pearlitic nose is laterally displaced to the right of the lower nose or region of bainitic transformation. For all of these alloys the end or tip of the pearlitic nose is somewhere between 1,200 and l,350 F. In other words, by the proper chemistry selection, i.e. balancing isothermal transformation retarding additions, and hardenability additions, it is possible to select an alloy possessing sufficient transformation time delays to permit working and cooling from the austenitizing temperature past the pearlitic nose, while permitting the bainitic transformation.

Returning now to the chemistry of the exemplary alloys of Table I, it will be noted that four elements (C, Mn, Mo and Cr) are conspicuous by their presence in all but two alloys. All of said elements are present in the alloys for their effects on hardenability, and the overall effect in retarding isothermal transformation. The hardenability of a ferrous alloy is governed in large part by its chemical composition, and each addition or element present in the alloy affects hardenability to some degree. Much of this is known so it will suffice to say that the degree differs between additions, and between varying sub-ranges within abroad range of elemental addition.

In the present invention, carbon, while inherently present in steel, should be present in an amount sufficient to yield precipitated carbides upon final cooling. Generally, it is desirable to include from 0.10 to 0.50 percent, by weight, in the alloy, but a broader range of 0.03 to 0.65 percent, by weight, may be used with success. As stated previously, all elemental additions affect hardenability. This may be extended to say that all elemental additions, to a large or small degree, are effective in retarding isothermal transformation. For instance, a number of elemental additions, such as M0, Cr, V, Si, Ti and Cb (identifiedfor convenience as Group I elements), are effective in pushing the pearlite nose to the right. One of the most effective of these additions is molybdenum. However, the' latter additions should be distinguished from some other elemental additions, such as Mn, Ni, B, N and Co, which act to shift the entire S curve to the right; these may be identified as Group II elements. Thus, while one or more Group I additions are desirable, it is preferable that the alloy contain at least 0.25 percent by weight, molybdenum. But, when a lesser amount is used, it is desirable to use at least 0.75 percent, preferably, 1.00 percent, by weight, chromium. These elements (Cr and Mo) and the others of Group I have been observed and are classified as ferrite stabilizers, and as noted above, all of said Group I elements retard the pearlite transformation, i.e., move the pearlite nose further to the right.

Considering now the procedural steps of this invention, it begins by heating a bainitic ferrous alloy to an austenitizing temperature of about l,500 to 2,200 F. where the first working, such as rolling, should start. Several reductions or passes, preferably at least four, on the order of about 10 percent minimum for each reduction, should occur between the starting temperature and the onset of the bainitic transformation or the B., temperature which occurs in the range of 900 to 1, 100 F. This sequence causes grain refinement of the austenite by repeated deformation and recrystallization. Additions of carbide forming elements may aid in the refinement of austenite grain size by precipitation in the austenite. At the lower temperatures, the recrystallization of austenite is retarded and the deformation causes substructural and textural development in the prior austenite grains. The further and final aspect of this procedure involves at least one additional pass I below the bainitic start temperature to provide additional dislocations which act as nucleation sites for carbide precipitation. The latter is necessary to form numerous, fine precipitates which increase strength and gated grain structure, with a lath-like substructure and random as measured by the nine-plane, (111) pole figure technique. A value of 1.0 represents a completely random orientation, whereas 9.0 shows 100 percent, or the orientation of a single crystal.

In contrast to this, the same alloy hot-rolled and finished above about 1,600 F. reveals a microstructure of largeequiaxedgrains (ASTM No. 3-5). While there is no lath-like substructure, large carbides, on the order of l to 5 pm in size, are found within the grains .and at the grain boundaries. This texture will nearly approach random with the intensity of any particular orientation below 2.5 times random, by the system noted above.

The improved properties of the alloys of this invention can best be demonstrated with rolling data and specific properties of six bainitic alloys treated by the method herein, and by the conventional hot rolling process, each of which are schematically illustrated in the FIGURE. For convenience, the chemistry, by weight percent, for the bainitic ferrous alloys are given in Table II, with the rolling data and property results in TABLE lV-continued Pass Sequence Thickness (inches) Reduction (7:)

*average elapsed time for alloys finished above 1590 F. was approximately 3 minutes; the elapsed time for alloys finished between 1 190-700" F. ranged between 10 and 26 minutes.

While temperatures were not ascertained following each pass for each alloy, where measured they were found to be linear, following a relatively constant cooling rate from the designated start to finishing temperature. For the alloys of this invention, the cooling rate varied between about -50 F ./min. The foregoing is merely representative of reduction conditions, as variations from this may be made, particularly with changes in alloy composition. However, under all reduction T bl 111, practices, it is essential that at least the final reduction TABLE II Alloy C Mn P 5 Si Ni Cr Mo V Ti B AC .46 .54 016 .014 .15 1.86 .87 .26 .005 .003 AD .41 1.36 015 .013 .33 .27 1.0 .46 .007 .001 A15 .17 .59 .012 .012 .34 .02 .01 .52 .002 .001 .003 AF .17 .50 012 .011 .25 .03 .01 .28 .001 .001 .003

TABLE lIl takes place within the bainitic region and concluded p ie before the complete transformation of the alloy to bai- Rollmg Temp. (F.) Transmon Temp. 35 Alloy sum Finish Y.S. (ksi) (Cv15 F) A l I600 650 1,1 1 80 Returning to the results of Table III, 1t Wlll be evident 2 1 600 700 120 6 that in nearly all situations wherein the method of this AA3 1600 800 114 80 invention was followed, higher strength and improved Q's: 1 238 g; :23 impact properties were realized. Thus, with the proce- AB2 15m 800 192 2()O dure detailed herein, it is now possible to gain optimum A33 2140 1725 mechanical and impact properties in as-rolled steels AC1 2080 1600 108 +200" AC2 1800 990 [06 without resorting to costly heat treatments. AC3 1600 1 1 75 1 :5 its) The attainment of optimum properties is clearly dem- AC4 1610 80 1 ADI 2020 1590 9] +200, onstrated with an exemplary comparison of Alloys A83 ADZ 1790 995 104 35 and AB2. The yield strength was increased by at least A133 1610 64 percent while the transition temperature was low- AD4 1600 1000 109 75 o o AE] 2000 1600 45 ered from 75 F. to -20O F. Another slgnificant dem- 2E3 38 3; h 8: onstration is noted when Alloy AB2 is compared to 1600 1000 H3 Alloy ABl, the latter having been finish rolled after AF] 2070 1600 46 75: complete transformation to bainlte. While the precise 2E :28 328,, critical temperatures of the worked alloy are not Am 1000 101 known, such an alloy under isothermal transformation To effect a proper comparison between the processed steel alloys of this invention and the conventionally rolled steel alloys, all reductions from 4 inches to /2 ;inch followed substantially one of two schedules set forth in Table lV.

conditions would reveal critical temperatures for the B and B, of approximately 920 F. and 750 F respectively. The critical temperatures for the worked alloy would be slightly higher. Nevertheless, it is quite clear that even 750 F., the isothermal B temperature, is well above the finish temperature of 620 F. for Alloy ABl. While the Y.S. thereof remained at about the same level as the Y.S. for Alloy AB2, it will be observed that a significant loss occurred in the transition temperature. That is, the transition temperature was raised from 200 F. to 50 F.

While it can be expected that improvement will vary with the specific alloy being treated, it is believed significant that large improvements in strength and/or impact properties were noted in all situations.

We claim:

2. The alloy according to claim 1 wherein the carbides range in size between about 0.01 to 0.03 pm.

3. The alloy according to claim 1 wherein the carbon is present in an amount between about 0.10 to 0.50

percent.

4. The alloy according to claim 1 wherein the intensity of said crystallographic texture ranges from 2.5 to 3.5 times random as measured by the nine-plane, inverse pole figure technique on the plane perpendicular to the working direction of said alloy. 

1. AN AS-WORKED BAINITIC FERROUS ALLOY CONSISTING ESSENTIALLY, BY WEIGHT, OF ABOUT 0.03 TO 0.65 PERCENT CARBON, A MINIMUM OF ABOUT 0.25 PERCENT MOLYBEDNUM, AN ADDITION OF AT LEAST ONE ELEMENT SELECTED FROM THE GROUP CONSISTING OF BORON, MANGANESE, NICKEL AND CHROMIUM, WITH THE BALANCE ESSENTIALLY IRON, CHARACTERIZED BY A CRYSTALLOGRAPHIC TEXTURE DOMINATED BY A (III) GRAIN ORIENTATION AND A MICROSTRUCTURE OF FINE, ELONGATED GRAINS, WITH A LATH-LIKE SUBSTRUCTURE, WHERE SAID LATH-LIKE
 2. The alloy according to claim 1 wherein the carbides range in size between about 0.01 to 0.03 Mu m.
 3. The alloy according to claim 1 wherein the carbon is present in an amount between about 0.10 to 0.50 percent.
 4. The alloy according to claim 1 wherein the intensity of said crystallographic texture ranges from 2.5 to 3.5 times random as measured by the nine-plane, inverse pole figure technique on the plane perpendicular to the working direction of said alloy. 